The properties of intermetallic compounds are somewhere between those of metals and ceramics. They frequently exhibit interesting application properties of these two classes of materials, such as high hardness, heat resistance, corrosion resistance at high temperatures, superconductivity, permanent magnetism or response to magneto-optical factors. There exist numerous investigations and publications dealing with the possible industrial applicability of intermetallic phases, of which the following publications are named as being representative: Intermetallische Phasen, Autorenkollektiv (Intermetallic Phases, Author Collective), published by VEB Deutscher Verlag fuer Grundstoffindustrie, Leipzig (1977); E. Fitzer in: Warmfeste und korrosionbesstandige Sinterwerkstoffe (Heat- and Corrosion-Resistant Sintered Materials), Metallwerk Plansee, Springer Verlag, Vienna (1956) and Intermetallic Compounds, J. H. Westbrook, J. Wiley & Sons, New York (1967).
Examples of such industrially used intermetallic phases are permanent alloys based on the intermetallic compound SmCo.sub.5 (J. Appl. Phys. 38, 1001 (1967)or high-temperature materials from MoSi.sub.2 (Fitzer, loc. cit.).
In many cases, the efforts to produce intermetallic compounds were finally suspended, partly because the preparation of promising intermetallic compounds failed and partly because products with the desired properties could not be obtained in an industrially usable or reproducible manner because of appreciable difficulties in the production.
Special difficulties arise with the production of intermetallic phases with an incongruent melting point. Incongruently melting intermetallic phases which are industrially important are, for example, Er.sub.3 Co, ErNi.sub.2, Er.sub.3 Ni, Fe.sub.2 Tb, Fe.sub.2 Nd, Nd.sub.3 Ni, Nd.sub.2 Fe.sub.14 B, SmCo.sub.5, Sm.sub.2 Co.sub.7, Nb.sub.3 Ge, Nb.sub.3 Sn, Nb.sub.3 Ti, Ni.sub.3 Al, Ti.sub.3 Al and TiAl.
The difficulties of producing single-phase, incongruently melting intermetallic compounds are described below using the intermetallic phases Nd.sub.2 Fe.sub.14 B, Sm.sub.2 Fe.sub.17 and .gamma.-TiAl.
1. Nd.sub.2 Fe.sub.14 B
The European patent 0 101 552 and the U.S. Pat. No. 4,756,775, to name a few publications, describe NdFeB alloys with exceptional magnetic properties.
The industrially important production methods for NdFeB are:
1. the conventional melt metallurgy with ingot casting, PA1 2. the calciothermal co-reduction and PA1 3. the melt metallurgy with rapid solidification. PA1 1. The Curie point of the intermetallic compound Nd.sub.2 Fe.sub.14 B is only T.sub.c =596.degree. K.; at T=423.degree. K., a re-orientation of the aligned spin is observed. With that, the application temperature of ternary NdFeB alloys is limited to about 400.degree. K. PA1 2. The NdFeB magnets, which are available at the present time, have insufficient corrosion resistance which is attributed to the increased formation of an Nd-rich phase. PA1 1. The nano-crystalline alloys also have a deficient heat stability. PA1 2. The isotropic, nano-crystalline powders, because of their isotropy, have comparatively low magnetic properties. PA1 3. The production of nano-crystalline NdFeB alloy powders by melt spinning is a technologically expensive process. PA1 4. The nano-crystalline NdFeB alloys are unsuitable for sintered magnets produced by powder metallurgical means. PA1 as described in J. Less-Common Metals 25, 131 (1971). PA1 thermochemical process; PA1 vacuum induction melting; PA1 skull melting with vacuum arc, plasma arc or electron beam; PA1 rapid solidification from the melt; or PA1 calciothermal co-reduction. PA1 a) adjusting the exothermic reaction of the calciothermal reduction by adjusting the oxide content of the reaction mixture, the composition of which corresponds to the desired single-phase alloy, in such a way that the temperature condition T.sub.m &gt;T.sub.R .gtoreq.0.9 T.sub.m (in .degree. K.) is fulfilled, T.sub.m being the melting temperature of the intermetallic phase and T.sub.R the reaction temperature; PA1 b) use of a reaction mixture, the components of which, with the exception of calcium, have an average particle size of .ltoreq.75 .mu.m; and PA1 c) tempering the reaction product at the end of the exothermic reaction at a temperature, which is at least 0.7 times the melting temperature T.sub.m of the desired, single-phase alloy, measured in .degree. K., but is less than the melting temperature T.sub.m, during a period sufficient for the diffusion of the components.
The conventional melt metallurgy (1) comprises the calciothermal reduction of the NdFe.sub.3 or the fused mass electrolysis to produce Nd metal and the subsequent vacuum inductive melting of the desired NdFeB alloy (K. Ohashi in: Proc. of the Gorham Advanced Materials Institute, Seattle, Wash. (1990); Metal Powder Report, MPR Publishing Services Ltd., Bellstone Shrewbury, Shropshire SY1HU England, 42, 6, 438 (1987)).
For the co-reduction method (2) different oxides are reduced calciothermally and transformed in a diffusion process into an NdFeB alloy. During the reduction and diffusion process, CaO is formed which is separated subsequently by wet chemical means from the NdFeB alloy (Proc. of the 8th Intern. Workshop REPM and their Appl., Dayton, 587 (1985) and Metal Powd. Rept. 42,438 (1987)).
The NdFeB alloys, produced by methods (1) and (2), are used to produce anisotropic sintered magnets by powder metallurgical methods.
The rapid solidification method (3) also starts with the vacuum-inductive melting of a preliminary NdFeB alloy which is then spun by melt spinning (J. Appl. Phys. 53, 2078 (1984)). A nano-crystalline NdFeB alloy with a crystallite size of &lt;400 nm is formed. The nano-crystalline NdFeB alloys have isotropic magnetic properties.
The NdFeB alloys, which are produced by melt metallurgical or co-reduction means and are obtainable commercially, usually have Nd or rare earth metal concentrations of 32 to 36% by weight.
In investigations of the microstructure of magnets of the composition Nd.sub.15 Fe.sub.77 B.sub.8, Nd.sub.14.5 Fe.sub.77.5 B.sub.8 and Nd.sub.16.7 Fe.sub.75.5 B.sub.7.8, the Nd.sub.2 F.sub.14 B and Nd.sub.1.1 Fe.sub.4 B.sub.4 phases, an Nd-rich phase and .alpha.-Fe were detected correspondingly.
The rapidly solidified, nano-crystalline NdFeB alloys usually contain Nd concentrations between 28 and 30% by weight. The spun NdFeB alloys are composed essentially of a two-phase structure of Nd.sub.2 Fe.sub.14 B and an Nd-rich grain-boundary compound. Only small proportions of a compound of the composition Nd.sub.2 FeB.sub.3 are present.
A synopsis of the permanent magnet alloys of today and of their applications is given in Proc. of the IEEE, 78, 6, 923 (1990). According to this, the available NdFeB alloys for sintered magnets have a remanence of B.sub.r =10 to 12.8 kG, a coercivity of iH.sub.c =8 to 24 kOe and energy products of (BH).sub.max =25 to 40 MGOe. The isotropic, spun NdFeB alloy powders have a remanence of only 8 to 9 kG, a coercive field strength of iH.sub.c =15 to 17 kOe and an energy product of (BH).sub.max =12 to 14 MGOe. The hot-molded, anisotropic, nano-crystalline alloy powders have clearly better magnetic properties. The remanence is B.sub.r =10 to 12.2 kG, the coercive field strength iH.sub.c =12 to 20 kOe and the energy product (BH).sub.max =22 to 35 MGOe.
The essential disadvantages of melt-metallurgical or co-reduced alloys for sintered NdFeB magnets are:
Magnets of nano-crystalline NdFeB powders have a lesser proportion of the Nd-rich phase. This brings about a corrosion behavior which is better than that of conventional sintered magnets.
The rapidly quenched NdFeB powders have the following summarized disadvantages:
On the other hand, of all the NdFeB alloys, the single-phase, incongruently melting, intermetallic compound Nd.sub.2 Fe.sub.14 B has, by far, the highest saturation magnetization and the best possible corrosion behavior.
Koon et al. (J. Appl. Phys. 57, 1, 4091 (1985) investigated the magnetic properties of R.sub.2 Fe.sub.14 B single crystals with R=Y, Nd and Tb. The single crystals were produced by means of the Czochralski method (B. N. Das, N. C. Koon: High Performance Permanent Magnet Materials, Proceedings of the Materials Research Society, Anaheim, USA, 96, 41 (1987)).
According to these investigations, Nd.sub.2 Fe.sub.14 B has a saturation magnetization at room temperature of 4.pi.M.sub.s =16.2 kG, which corresponds to a maximum, theoretically attainable energy product of (BH).sub.max =65.6 MGOe.
However, it was shown in numerous investigations that single phase Nd.sub.2 Fe.sub.14 B does not have any outwardly measurable coercive field strength and only an inadequate sintering behavior. According to these investigations, a non-magnetic phase, such as the Nd-rich phase, is required in the multi-phase NdFeB sintered magnets on one hand, in order to produce a macroscopically measurable coercivity and, on the other, to promote compacting during the manufacture of the magnet by the formation of a liquid phase during the sintering process.
The European patent 0 249 973 discloses that an improvement in the magnetic properties and the corrosion resistance is possible by the separate production of Nd.sub.2 Fe.sub.14 B and a non-magnetic sintering aid. To begin with, an NdFeB alloy having the composition Nd.sub.13 Fe.sub.81 B.sub.6 and various sintering aids are melted separately by vacuum induction, milled to an average particle size of 3 .mu.m, mixed in a ball mill, hot pressed isostatically and treated with heat. The two-phase magnet, with the composition of (Nd.sub.13 F.sub.81 B.sub.6).sub.96 Al.sub.4 with B.sub.r =14.9 kG, iH.sub.c =10.1 kOe and (BH).sub.max =49.7 MGOe, has the most advantageous magnetic properties of all the magnets investigated.
The corrosion behavior of different two-phase magnets was investigated in different corrosion tests in the coated and non-coated state and compared with an Nd.sub.14 Fe.sub.80 B.sub.6 sintered magnet which had been produced by conventional powder metallurgical means. For example, the NdFeB magnets were exposed to a corrosive atmosphere with a moisture content of 90% at 60.degree. C. for a period of 100 hours. The conventionally produced Nd.sub.14 Fe.sub.80 B.sub.6 magnet showed signs of corrosive attack already after one hour. On the other hand, the unfavorable two-phase (Nd.sub.2 Fe.sub.14 B).sub.89.5 (NdCu.sub.2).sub.10.5 magnet already showed a 10-fold and the most stable (Nd.sub.2 Fe.sub.14 B).sub.95 (Nd.sub.30 Fe.sub.40 Al.sub.30).sub.5 a 50-fold higher corrosion resistance.
The compositions of the different structure components were determined by micro-probe analyses. According to these, the hard magnetic Nd.sub.2 Fe.sub.14 B phase (.PHI. phase) has an average composition of 12.5 at % Nd, 81.5 at % Fe and 6 at % B.
According to investigations by Durst et al. (Proc. of the 5th Intern. Symp. of Magn. Anisotropy and Coercivity in RE-Transition Metal Alloys, Bad Soden, FRG, 2, 209 (1987)), a higher temperature stability due to an improved structure morphology is achieved by the two-phase structure of the NdFeB magnet, consisting of the hard magnetic .PHI. phase and a non-magnetic, Nd-rich phase.
Nd.sub.2 Fe.sub.14 B (.PHI. phase) is an incongruently melting intermetallic compound with a melting point of 1180.degree. C. At the peritectic melting point, the .PHI. phase decomposes into an Nd-rich melt and .alpha.-Fe which leads to a drastic decrease in the coercive field strength and is therefore to be avoided. Chin et al. (Proc. of the loth Intern. Workshop on Rare Earth Magnets, Kyoto, Japan, 2, 451 (1989)) investigated the .PHI. phase for a possible expansion region. According to this investigation, the .PHI. phase at 750.degree. C. has only one defined composition, but at 1090.degree. C. and 1135.degree. C. a homogeneity range of 0.3 at %.
The stoichiometric .PHI. phase has a composition of 11.8 at % Nd, 82.3 at % Fe and 5.9 at % B.
This leads to the conclusion that the composition, given in the European publication 0 249 973 for the 2-14-1 phase, cannot correspond to the single-phased .PHI.-phase and that proportions of extraneous phases must be present.
For incongruently melting intermetallic phases, production methods based on solid reactions are to be preferred in order to avoid the incomplete peritectic formation of the desired phase which is unavoidable when cooling from the melt.
A manufacturing process for nano-crystalline alloys in the solid phase, which has been known for many years, involves the mechanical alloying of metallic alloy components. Schultz et al. (Proc. of the 5th Intern. Symp. on Magn. Anisotropy and Coercivitity in RE-Transition Metal Alloys, Bad Soden, FRG, 1, 301 (1987) investigated the formation of NdFeB by mechanically alloying the elementary powders and determined the magnetic properties of the nano-crystalline alloying powders prepared. It was observed that there was no formation of the intermetallic Nd.sub.2 Fe.sub.14 B phase during the mechanical alloying. Instead, an amorphous mixture of Nd and B is formed as well as crystalline .alpha.-Fe which is alloyed with 4 to 5 at % of Nd. The hard magnetic Nd.sub.2 Fe.sub.14 B phase is formed only in a subsequent thermal treatment by solid reaction of the nano-crystalline elementary powder mixture. An isotropic, nano-crystalline NdFeB powder is formed with magnetic properties comparable to those of the known, rapidly chilled alloy powders.
An economically and industrially accessible method for the concerted, exclusive preparation of the intermetallic Nd.sub.2 Fe.sub.14 .sup.B phase is thus not known.
2. Sm.sub.2 Fe.sub.17
The intermetallic Sm.sub.2 Fe.sub.17 and the compound Sm.sub.2 Fe.sub.17 N.sub.3 obtained from it by nitration have metallurgical and magnetic properties similar to those of the Nd.sub.2 Fe.sub.14 B phase.
Knowing the intermetallic Nd.sub.2 Fe.sub.14 B phase, intensified research was conducted for new magnetic materials. Emphasis was placed on the different RE.sub.2 TM.sub.17 -intermetallic phases and their nitrogen compounds of the RE.sub.2 TM.sub.17 N.sub.x type, where RE comprises all rare earth metals including yttrium and TM comprises the metals Fe, Ni or Co. Sm.sub.2 Fe.sub.17 N.sub.3 has proven to be particularly promising. For the preparation of the nitride, the synthesis of a single-phase Sm.sub.2 Fe.sub.17, which was subsequently nitrated at 400.degree. to 500.degree. C., was attempted first.
The objects of the European publication 0 369 097 are, among others, nitrides of the composition RE.sub..alpha. Fe.sub.(100-.alpha.-.beta.-.gamma.) N.sub..beta. H.sub..gamma. with .alpha.=5 to 20 at %, .beta.=5 to 30 at % and .gamma.=0.01 to 10 at %.
Sm.sub.2 Fe.sub.17 is an incongruently melting intermetallic phase with a melting point of 1280.degree. C. At the melting point of Sm.sub.2 Fe.sub.17, there exists the thermodynamic equilibrium: EQU Sm.sub.2 Fe.sub.17 =.alpha. Fe+Sm-rich melt
In investigating the magnetic properties of Sm.sub.2 Fe.sub.17 N.sub.3-, it was noted by means of the hysteresis curve that residues of .alpha.-Fe are still present. Such presence should be prevented if at all possible.
In order to compensate for the precipitation of .alpha.-Fe, which is unavoidable during the melt metallurgical preparation of Sm.sub.2 Fe.sub.17, diffusion annealing is required.
The Curie temperature of the 2/17-phase is T.sub.c =130.degree. C. By nitrating Sm.sub.2 Fe.sub.17 and forming Sm.sub.2 Fe.sub.17 N.sub.3-.gamma., the Curie temperature is increased to T.sub.c =470.degree. C. The saturation magnetization of the nitride was determined to be 4.pi.M.sub.s =1.54 T. This corresponds to a theoretically attainable energy product of (BH).sub.max of 59.3 MGOe.
This compound also, like the single-phase Nd.sub.2 Fe.sub.14 B, has practically no coercivity. As in the case of Nd.sub.2 Fe.sub.14 B, a coercivity can be produced by separating the grains by an intercrystalline, non-magnetic second phase similar to the Nd-rich phase.
A second possibility for obtaining a coercivity is by setting a nano-crystalline structure. Schnitzke et al. (K. Schnitzke, L. Schultz, J. Wecker, M. Katter, submitted to Appl. Phys. Lett.) initially synthesized a nano-crystalline Sm.sub.2 Fe.sub.17 by mechanically alloying the elementary powder. Radiographic analyses of the powder revealed only an amorphous phase and crystalline .alpha.-Fe. The intermetallic Sm.sub.2 Fe.sub.17 phase is formed only during a subsequent heat treatment and was then converted by nitration to the Sm.sub.2 Fe.sub.17 N.sub.3-.gamma. compound.
It is a disadvantage of these synthesis methods that only isotropic Sm.sub.2 Fe.sub.17 N.sub.3-.gamma. alloys can be produced. Industrial production aiming at single phase Sm.sub.2 Fe.sub.17 is thus not yet possible by economic means.
3. .gamma.-TiAl
The areas of turbine manufacture, aircraft manufacture and space technology show a precipitous increase in material research in the field of intermetallic compounds. Alloys based on nickel find the most frequent application. However, the high densities of these alloys prove to be disadvantageous, for instance, in view of the centrifugal forces that arise in the turbines.
Aside from a high hardness and thermal stability, the intermetallic phases usually have a comparatively low density. This is so particularly for the aluminides of titanium. Alloys based on Ti.sub.3 Al (M R S Symp. Proc., Boston, 39, 221 (1985) and Z. Metallkde. 80, 5, 337 (1989)) were the first to be used industrially. However, because of a deficient oxidation resistance, their use is limited to a maximum temperature of 973.degree. K. This contributed to aluminides of titanium with a higher aluminum content being investigated more thoroughly from the point of view of increasing the use temperature. The focus of attention of this research was directed on the aluminide .gamma.-TiAl.
.gamma.-TiAl solidifies through a peritectic reaction, has a melting point of 1480.degree. C. and a homogeneity range of approximately 49 to 56 at % Al at room temperature (A. Hellwig, Dissertation of the University of Dortmund, Dortmund, 1990) and can be synthesized by the following methods:
Further details may be found in JOM, 42, 3, 26 (1990); JOM, 43, 5, 30 (1991); JOM, 42, 3, 16 (1990); Mat. Sci and Eng., 269 (1988) and JOM, 42, 3, 22, 1990).
Skull melting and the rapid solidification from the melt by, for example, melt spinning, have gained special importance for the melting of titanium alloys and aluminides of titanium.
Skull melting in the vacuum arc is a method, which has long been known and accepted as comparatively simple technologically and inexpensive. .gamma.-TiAl solidifies incongruently so that conventional melt metallurgical methods, such as skull melting, do not ensure uniform, chemical homogeneity. By rapidly chilling from the melt, a distinctly better homogeneity of the .gamma.-TiAl and a significantly finer grain are produced which increase the ductility at room temperature appreciably.
It is a disadvantage of the rapid chilling of TiAl from the melt that this method is more complicated technically and therefore more expensive.
Moreover, the metallic starting components of the alloy or the intermetallic phase are required for all melt metallurgical processes. Melting, moreover, is appreciably more difficult for oxygen- and/or nitrogen-affine metals, such as Ti and Al.
In contrast to melt metallurgy, the more stable and more inexpensive oxides TiO.sub.2 and Al.sub.2 O.sub.3 can be used as raw materials for the calciothermal co-reduction.
The European patent 0 039 0791 discloses that Ti-Al alloys with a basic composition corresponding to TiAl.sub.6 V.sub.4 (based on % by weight) and different other additives can be produced by calciothermal co-reduction. An alloy powder with a homogeneous grain structure of .alpha.- and .beta.-Ti crystals is formed.
In the Proc. of the 6th World Conference on Titanium, Cannes, France, 2, 895 (1988), the calciothermal preparation of Ti-Al alloys is described and the preparation of a single phase .gamma.-TiAl is reported. As raw material mixture, the mechanically mixed oxides or hydroxides or the calcined oxides were used. The latter were prepared by precipitation from an HCl or H.sub.2 SO.sub.4 solution, annealing in air and grinding to fine powder with a particle diameter of 10 .mu.m. The reducing temperature of the reaction mixture should be between 1173.degree. and 1273.degree. K. and the reducing time should be 19 to about 36 hours.
There is, however, considerable doubt as to whether the method described actually results in the formation of single-phase intermetallic alloys of the composition .gamma.-TiA1. For example, it is known that the reduction reaction between TiO.sub.2, Al.sub.2 O.sub.3 and Ca is extremely exothermic and spontaneous. Therefore, when the aforementioned oxides are used exclusively, a reducing temperature of 1173.degree. to 1273.degree. K. and reducing times of 1 to 36 hours are improbable. Presumably, the "reducing temperature" is not meant to be the reaction temperature during the exothermic reaction but the furnace temperature. For an incongruently melting, intermetallic compound, it is not sufficient to state the furnace temperature, because the temperature must lie below the melting point of the intermetallic phase at every point in time of the process. The reaction temperature during the exothermic reducing reaction therefore also has a particular importance.
The use of hydroxides as raw materials is likewise improbable. Hydroxides decompose during the reduction process to an oxide and H.sub.2 O. The latter reacts further with Ca to form H.sub.2 and CaO which can lead to an oxyhydrogen gas reaction and, from a safety point of view, makes the implementation of the process with hydroxides as raw materials implausible. Accordingly, this reference is also eliminated as a source that teaches the synthesis of single-phase, incongruently melting compounds of the .gamma.-TiAl type.